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In situ investigation of MgO nanocube deformation at
room temperature
Inas Issa, Jonathan Amodeo, Julien Réthoré, Lucile Joly-Pottuz, Claude
Esnouf, Julien Morthomas, Michel Perez, Jérome Chevalier, Karine
Masenelli-Varlot
To cite this version:
In situ investigation of MgO nanocube deformation
at room temperature
I. Issa
1,2, J. Amodeo
1*, J. Réthoré
2, L. Joly-Pottuz
1, C. Esnouf
1, J. Morthomas
1,
M. Perez
1, J. Chevalier
1and K. Masenelli-Varlot
11 MATEIS, CNRS UMR5510, Université de Lyon, INSA-Lyon, F-69621 Villeurbanne Cedex, France
2 LAMCOS, CNRS UMR5259, Université de Lyon, INSA-Lyon, F-69621 Villeurbanne Cedex, France
(*) Corresponding author:
Dr. Jonathan Amodeo
Laboratoire MATEIS - UMR CNRS 5510 Groupe METAL, groupe CERA
25 avenue Jean Capelle Bat. St. Exupéry, 2ème étage 69621 Villeurbanne Cedex, France Tel: 04 72 43 82 35
Fax: 04 72 43 85 39
In situ investigation of MgO nanocube deformation
at room temperature
Abstract
The mechanical behaviour of <100>-oriented MgO nanocubes is investigated using in situ TEM compression tests at room temperature and molecular dynamics simulations. Experiments show high strength and ductility in addition to specific deformation mechanisms interpreted by the simulation. The nucleation and the propagation of ½<110>{110} dislocations are at the onset of the plastic deformation. The different deformation processes as well as the possible formation of a dislocation network during compression are discussed.
Keywords: MgO, in situ TEM, nanoparticles, molecular dynamics, nanomechanics, dislocations
I
Introduction
strength and ductility [5]. Owing to the characteristic sizes of such objects, in situ mechanical tests in a transmission electron microscope (TEM) or in a scanning electron microscope (SEM) have revealed to give valuable pieces of information regarding the deformation mechanisms [6,7]. Deep in the submicrometer scale, size phenomena influence elastic properties of nanowires (NWs) and nanoparticles through the combination of surface and core effects in both the experiments and atomistic simulations [8,9]. This size effect involves also the plastic deformation regime for whom few theories based on the single-arm dislocation source model or the dislocation nucleation/starvation model [10,11] still try to provide a unique description of the now well-known principle “smaller is stronger”. Nevertheless, most of these studies are carried on metals, especially face-centered cubic (FCC), and only few works have been dedicated to ceramics [12-15]. Calvie et al. and collaborators have recently reported in situ compression tests in the TEM of a g-alumina Al2O3 nanospheres [14]. The main feature of this study is that the sample undergoes wide
and homogeneous plastic deformation, as observed in metallic [16,17], intermetallic [18] or silicon [19,20] nanoparticles. A detailed mechanical analysis based on digital image correlation (DIC) and finite elements simulations allowed the determination of a mechanical constitutive law [21]. However, the identification of elementary mechanisms responsible for plastic deformation in g-Al2O3 nanospheres could not be determined in situ. In particular, no
dislocation could be observed experimentally. This lack was primarily attributed to the spherical geometry of the sample and to the diffraction conditions. Furthermore, g-Al2O3 is
Here we propose to reach further the investigation of the mechanical behaviour of single crystalline ceramic nanoparticles using in situ TEM compression tests and molecular dynamics (MD) simulations applied to magnesium oxide (MgO).
In this study, we report in situ compression experiments in the TEM and MD compression simulations of MgO nanocubes at RT. Simulations are performed to investigate elementary deformation mechanisms at the atomic scale and corroborate experimental observations.
II Materials and methods
II.1 Nanocube experimental synthesis and characterization
Magnesium oxide nanocubes were prepared by burning commercial Magnesium chips (4-30 mesh) in air as described in the pioneer work of Heidenreich [44]. This specific method is known to produce perfectly cubic-shaped nanoparticles down to sizes of a few nm. The smoke particles were caught directly on three substrates: a glass substrate for SEM imaging, a TEM grid (Cu 300-mesh covered by a holey carbon film) and the nanocompression sapphire substrate.
Primary characterisations were carried out by SEM, high resolution TEM (HRTEM) and weak beam dark field (WBDF) TEM. SEM images were acquired on a Zeiss Supra 55VP microscope. The acceleration voltage was fixed to 1 kV and images were acquired with a secondary electron Everhardt-Thornley detector. For the HRTEM observations, a JEOL 2010F TEM microscope equipped with a field emission gun and operating at 200 kV was used. Images were recorded using a Gatan Orius 200 CCD camera. Energy-dispersive spectroscopy (EDS) was performed using an 80 mm2 SSD detector from Oxford Instruments.
Finally, WBDF characterisations were carried out on a JEOL 200CX microscope, equipped with a tungsten filament that operates at 200 kV. Images were acquired with the wave [220], that allows the detection of dislocations in the (101), (101&), (011) and (011&) slip planes. Under such conditions, dislocations can be imaged as narrow lines which are approximately 10-15 nm wide [45].
composed of MgO (99.9 ± 0.1 at.%). Particles are single crystals, as shown in the HRTEM image of figure 1b. The surfaces of the nanocubes are crystalline as well and are oriented along the <100> directions, as shown by the electronic diffraction patterns, figure 1b. Regarding the importance that dislocations have during all stages of plastic deformation, a large number of MgO nanocubes were characterized by WBDF before any mechanical test. Figure 1c displays a WBDF image of two MgO nanocubes. The image shows no bulk lattice defect such as dislocations, grain boundaries or twinning. Contrasts were only observed at the contact points of the nanocubes, and were attributed to the stress field induced by the elastic compression of the lattice [46]. None of the imaged nanocubes were found to contain bulk lattice defects. In the following, similar samples used for in situ nanocompression testing will thus be considered as initially dislocation-free.
II.2 In situ TEM nanocompression
Pristine MgO nanocubes were deposited onto a 75 µm-thick sapphire substrate by passing the substrate in the smoke during synthesis. This method precludes the use of solvents that may modify the surface structure of MgO samples [47]. Since the nanocubes exhibit <100> surfaces, the compression axis is expected to be parallel to [001].
In situ nanocompression tests were carried out using a dedicated sample holder from Nanofactory Instruments, fitted in a JEOL 2010F microscope operating at 200 kV accelerating voltage. The sample holder was equipped with a truncated diamond tip (flattened area of about 0.25 μm2) and a load cell (maximum load of 3 mN), as shown figure
1d. Particles were positioned on the substrate, displaced toward the tip during compression at a controlled displacement rate of 2 nm/s that is equivalent to an engineering strain rate of 0.02 s-1 for a 100 nm nanocube.
fields to be estimated with a sub-pixel resolution (about 1/100 pixel). The distance measured between substrate and diamond tip edges corresponds to the instant size of the sample. Furthermore, not only the longitudinal displacement but also the lateral deformation was obtained. As the longitudinal and lateral deformations were synchronized on-the-fly with the force value, true stress and true strain were deduced considering that lateral deformations were equivalent in both directions. The true strain is defined as ln (𝐿 𝐿⁄ *), where 𝐿 is the instantaneous position and 𝐿* is the initial position of the substrate with respect to the tip i.e., the initial nanocube size. The true stress is defined as the ratio between the instant force measured by the sensor and the effective cube surface inferred from the DIC calculation.
Fig. 1 (Color online) (a) SEM image of smoked MgO nanocubes. (b) HRTEM image of a
nanocompression simulation set-up, made of two force fields to respectively sustain and compress the nanocube.
II.3 Computational methods
MD simulations were performed to investigate elementary deformation mechanisms of initially dislocation-free MgO nanocubes under compression at RT using the LAMMPS code [48]. Atomic interactions were described using a rigid ion model that included a Buckingham term in addition to long-range Coulombic interactions. Here we used the Ball and Grimes partial charges parameterization [49] successfully employed to describe surface diffusion [49], elastic constants and dislocation properties [50], which are considered as key points for interatomic potential transferability towards nanomechanical test simulations. We used a cut-off parameter of 8 Å for short-range interactions. Full long-range interactions between charges in compression simulations were computed using the multi-level summation method (MSM) solver [51]. MSM relative error in per-atom forces from 10-4 to 10-8 were tested without
significant outcome on the simulation results. Cubic samples with edge lengths from 4.2 nm to 12.7 nm were shaped with free surfaces oriented along the <100> directions as suggested by the experiment. The MD compression tests of MgO nanocubes were performed using the following procedure. First, the structure of the nanocubes was optimized using conjugate gradient and the FIRE algorithm [52]. Next, the samples were equilibrated during 30 ps in the NVE ensemble down to 300 K. Then, we used the Nosé-Hoover thermostat [53] for 50 ps equilibration in the NVT canonical ensemble. Finally, the compression tests were performed using two external potentials (see figure 1e) which model an infinite flat punch and the substrate [54,55]. To model uniaxial compression along the [001] direction, the top indenter was subjected to a constant displacement rate equivalent to an engineering strain rate of 108
s-1. The bottom potential was kept fixed to sustain the sample. During compression
temperature of 300 K. Note that sample size, strain rate and other simulation parameters were chosen regarding settings generally adopted in the case of the embedded-atom method (EAM) for metals [17], what involves a significant increase of the cpu costs due to long-range Coulombic interactions in this study. The true strain was calculated as in the experiment (see previous section) and the compressive stress was computed as the ratio between the force experienced by the indenter and the instantaneous maximum contact area drawn by the last upper atomic layers. Simulations were analysed using AtomViewer [56], a tool which combines a modified bond-angle method [57] to identify crystalline structures, a Nye-tensor analysis for dislocation Burgers vector definition [58] and a skeletonization algorithm for dislocation reconstruction [59].
An Ewald summation method [60] with a radius of 12 Å was used for the calculation of bulk material properties that involved periodic boundary conditions e.g., lattice parameter, elastic constants and generalized stacking fault (GSF) energies [61]. We used the same short-range interaction cut-off (8 Å) for bulk material properties and nanocompression simulations. GSF energies were computed using the full periodic method described in Gouriet et al. (2014) study [62] to avoid artefacts induced by charged surfaces. Calculations were performed along the ½<110> Burgers vector direction for {110}, {100} and {111} planes. Simulation cells were chosen large enough, especially along the orientation of the stacking fault normal, to reduce the interactions between fault periodic replicas and minimize long-range Coulombic effects. The initial size of the simulation cells are 1.79 nm*1.69 nm*14.29 nm, 1.79 nm*1.79 nm*14.31 nm and 1.49 nm*1.55 nm*27.71 nm respectively for {110}, {100} and {111} GSF calculations.
III Results
III.1 Mechanical behaviour of MgO nanocubes
stress-strain curve, four on the fly on-the-fly images are presented figure 2b-e for true strain of about 𝜀=1.2%, 𝜀 =18.7%, 𝜀 =49.7% and 𝜀 =78.9% respectively. The sample is deformed up to 𝜀 =79.0% before final unload.
Early in the first cycle, the stress-strain dependency is nonlinear and noisy. This transitory stage reflects the accommodation regime between the substrate, the sample and the tip up to 𝜀 =3.5%. Then, the curve exhibits an elastic-like behaviour up to 𝜀 =9.0% and a true stress of about 𝜎=0.78 GPa. A Young’s modulus of about 𝐸[**0]=141.9 GPa is deduced from the early beginning of the first unload. From this point and during the rest of the whole compression test, mobile contrast bands are observed in the nanocube.
Fig. 2 (a): Stress-strain curve for a 140 nm edge lengths MgO nanocube compressed in
Contrast bands may occur for several reasons e.g., the elastic deformation of the lattice or the presence of mobile/sessile dislocations. In the following, we will focus only on contrast bands that are believed to be the signature of dislocations. These contrast bands may be recognized by their peculiar curved shape as well as their specific orientation e.g., the two angled contrast bands observed on the 𝜀 =18.7% image of figure 2. Nevertheless, due to the Bragg conditions, all the defects cannot be observed during the experiment.
One can note that all the residual contrasts accumulated during the first load cycle, up to 𝜀=13.1%, vanish during the first unload. This phenomenon occurs because of both the compressive stress relaxation and the surface image forces, believed to be particularly effective in nanometer-sized samples. This leads to a perfectly refreshed microstructure comparable to the one observed figure 2b.
cycle and a total deformation of about 𝜀=79.0%. No sign of failure has been observed during the unload.
III.2 Focus on the early stages of deformation
In order to investigate further the elementary mechanisms are responsible for the first stages of the MgO nanocubes deformation, we performed two supplementary in situ compression tests. Stress-strain curves of 90 and 120 nm edge lengths nanocubes are represented figure 3. As the stress-strain curve main features of ca. 100 nm range MgO nanocube have been described in the previous section, we will focus only on new outcomes in the following. We observe that both samples are characterized by a strain burst occurring at higher stress than in the case of the 140 nm sample. At this point, true stress is about 2.72 and 5.03 GPa, respectively for the 120 and the 90 nm samples. In the inset of figure 3, an image of the 120 nm sized nanocube is shown at 𝜀=7.3% i.e., during the initial strain burst. Once again, we observe an inclined and straight contrast band, tilted of about 45 degrees from the indenter and generated from the surface and/or an edge of the nanocube, similarly to what was observed in the case of the 140 nm sample (figure 2).
Fig. 3 (Color online) Stress-strain curves for two MgO nanocubes of 90 (red line) and 120
compression, 𝜀=7.3%. The black arrow shows a contrast band corresponding to dislocations that emerge from the surface of the cube. Two MD simulation snapshots of a 12.6 nm edge lengths nanocube are also represented. Reconstructed ½<110>{110} dislocations are in green. The slip plan is in red. Atoms are shown in transparent light grey for the sake of clarity. The coordinates system is oriented along the cubic directions.
simulated stress-strain curves by the correlated stress drops. When dislocations nucleate (or escape), top or bottom surfaces of the sample are rearranged leading to a slight increase of the distance between the sample top surface and the indenter. As the flat punch force varies inversely with this distance, the stress drops down. The onset of plastic deformation is controlled by ½<110>{110} dislocations only, nucleated either from an edge, a surface or a corner of the nanocube. A detailed analysis of the localization of dislocation nucleation first events is summarized table 1.
Tab. 1 Young’s moduli and yield stresses for MD compression simulations, plus sites and
slip systems of the first nucleated dislocations. 𝐸2345=192.3 GPa is deduced from the 𝐶9: and the anisotropic elastic theory. For the 7.6 nm particle, two dislocations nucleate simultaneously.
Fig. 4 (Color online) Stress-strain curves of MgO nanocubes from MD compression
simulations. - : Images of nanocubes during compression. Green lines correspond to dislocations and red arrows represent ½ <110> Burgers vector orientation. - show the evolution of the 5.9 nm sample. - show the dislocation organization at the end of the first nucleation peak of the 12.6 nm sample. The blue circle shows a dislocation junction embryo. The coordinates system is oriented along the cubic directions.
IV Discussion
IV.1 Toward small-size effects on MgO mechanical properties
On the contrary to their bulk counterpart, nanometer-sized MgO particles deform up to large strain, beyond 𝜀=78.9% in the experiment, and no sign of crack has been observed during and after the compression of the nanocubes. These results apply also in the case of MD compression simulation using the Ball and Grimes interatomic potential. In the simulation, we observe a strong effect of the size decrease on the elastic regime. For particles with size lower than 10 nm edge length, 𝐸[**0] decreases reducing size (table 1). Furthermore, the elastic regime becomes non-linear. This phenomenon is weaker for larger particles and is thus not expected in the experiments (at least for the sizes we investigate here). Nevertheless, 𝐸[**0] is still lower in the experiment (~142 GPa for the 140 nm particle) compared to the experimental bulk value 𝐸[**0]2345~248 GPa [71,72]. We believe that this variation is not a size-effect but may rather be an extrinsic effect as e.g., electron-beam assisted deformation [73-75]. In Zheng et al. (2010), the authors point out that the force required to deform amorphous silica (a-SiO2) nanoparticles at a given elastic strain is
lowered by a factor 2 to 3 in the case of on tests compared to electron-beam-off [74]. In Mačković and collaborators (2014), a-SiO2 nanoparticles are pre-irradiated and
then compressed under beam-on/off conditions. Results show that the Young’s modulus is increased by a factor 2 (whatever the conditions) compared to the low-dose/beam-off reference conditions with a maximum load force of 50% up to 500% the original one [73]. Similar observations have been made by Zhang and collaborators in crystalline zinc tin oxide NWs where electron-beam irradiation is believed to change elastic and electrical conductivity properties [75]. Nevertheless, despite the fact that similar processes might influence the Young’s modulus in our experiments, we believe that they should not modify the elementary mechanisms responsible for plastic deformation.
size variation, during the deformation of the 140 nm sample, on the subsequent elastic reloads. As a corollary, no effect of the beam exposure time has been deduced from the analysis of elastic reloads. The aspect size effect on the subsequent elastic portions observed in the simulations has also been investigated without significant outcome. More details about this analysis are provided as supplementary information. Finally, the flow stress is also influenced by downscaling, and raises up in comparison to bulk <100>-compression tests at RT i.e., from ca. 50 MPa in bulk conditions [24,26,27,76] up to the GPa range for the nanocubes (figures 2 and 3). These results confirm that lowering the scale permits to increase both strength and ductility even for originally brittle ceramics. More specifically, yield stresses from experiments and simulations seem to exhibit a size-effect around 10 and 100 nm. However, the number of experiments should be increased, as the investigated range of size, to get a more significant statistic and further strengthen this hypothesis. In the following, we will further focus on the special features of the plastic deformation regime of MgO nanoparticles.
IV.2 Dislocation nucleation and slip systems in nanosized MgO
charge repulsion between ions [80]. This phenomena leads to a surprisingly high shear stress for the {111} slip planes compared to what is generally observed e.g., in FCC metals. With this orientation, {100} slip can only occur in case of misorientations or local rotations. Nevertheless, ½<110>{100} slip is believed to be effective under a rather higher shear stress than ½<110>{110} in RT compression experiments [25,27,79], what does not make {100} slip planes suitable candidates to accommodate deformation in our study.
Although the yield strength in the MD simulations (figure 4) is obviously influenced by the strain rate dependence of dislocation nucleation from the surfaces [67], the role of ½<110>{110} dislocations during the compression of MgO nanocubes is most probably strain rate independent and further constrained by energetic considerations.
To further investigate the relative role of {100}, {110} and {111} glide planes, we have calculated GSF energies using the same interatomic potential than for the MD compression simulations. GSF energies are computed by simply shifting the atoms contained in the upper half of a simulation cell relatively to its lower half by an appropriate translation vector owned by the boundary plan. GSF energies provide a good estimate of the sensitivity of a slip plane to be sheared in a given direction. This concept is frequently used to discuss dislocation-based elementary processes as dislocation core spreading through Peierls-Nabarro approaches [61,81,82] or dislocation nucleation [83-85]. As shown in e.g., Carrez et al. (2009) [86], GSF energies calculated in the {110} and the {100} planes of MgO do not exhibit the characteristic, “FCC-like”, stable stacking fault energy (sSFE) what allows ½<110> undissociated dislocations only. This corroborates the results inferred from the Nye tensor analysis during the MD compression simulations where perfect dislocations only are observed (figure 3 and figure 4). Conversely, unstable stacking fault energies (uSFE) can be computed along the Burgers vector direction. The uSFE increases up to a maximum value 𝑢𝑆𝐹𝐸ABC reached for a displacement equivalent to the half of the Burgers vector length (i.e., 1.49 Å). Here we find 𝑢𝑆𝐹𝐸{00*}ABC =916.7 mJ/m2, 𝑢𝑆𝐹𝐸
{0**}ABC =2073.1 mJ/m2 and 𝑢𝑆𝐹𝐸{000}ABC=2309.4 mJ/m2. 𝑢𝑆𝐹𝐸
and recent molecular statics simulations using the same interatomic potential [50,86]. While the difference between sSFE and uSFE is often used as an energetic criterion to justify preferred dislocation-based nucleation processes (i.e., partial dislocation, perfect dislocation or twinning) in FCC metallic nano-objects [83], we believe that in the strict case of perfect dislocation, the height of the energetic barrier to produce an elementary shear equal to the Burgers vector (i.e., the 𝑢𝑆𝐹𝐸ABC) is a good estimate of the slip plane sensitivity to dislocation nucleation. Here we show that 𝑢𝑆𝐹𝐸{00*}ABC < 𝑢𝑆𝐹𝐸
{0**}ABC < 𝑢𝑆𝐹𝐸{000}ABC, what confirms and the occurrence of ½<110>{110} dislocations, and the lack of ½<110>{111} dislocations in the MD compression simulations of MgO nanocubes and thus possibly in the in situ TEM experiments. Finally, one can note that 𝑢𝑆𝐹𝐸ABC values are one to two orders of magnitude higher in MgO than the sSFE-uSFE range commonly reached in standard FCC metals i.e., on the range of 10-100 mJ/m2 [87-89], what might support the extremely high strength
observed on the simulation simulated stress-strain curves (figure 4).
IV.3 Deformation regimes and dislocation network
that induce localized deformation and directly escape from the cube after their nucleation, as observed in figure 4. This dislocation starvation process is less pronounced in the 7.9 nm edge length sample and disappears nearly from the 12.6 nm sample. Actually, in the case of the 12.6 nm sample, deformation is homogeneous and dislocations nucleate from multiple slip systems during the first stress drop. In this case, gliding dislocations intersect and react what can be viewed as the critical step of the dislocation network formation (figure 4). This process is strongly different from the dislocation starvation process described above, as it does not require further dislocation nucleation i.e., dislocations later unlock from the dislocation network and multiply to accommodate further deformation.
Fig. 5 (Color online) TEM images of a 140 nm MgO nanocube during the first cycle of load.
One should note that the lack of subsequent large stress drops on the 12.6 nm sample stress-strain curve is not only due to this last microstructural process but also to the increase of the sample size. Indeed, following a simple first order approach, one can approximate the average shear 𝛾 produced by a {110} gliding dislocation in a cube of edge length 𝑙 by equation (1):
𝛾 = 𝑏I J= 𝑏
√7
74 (1)
Where 𝑏 is the Burgers vector length, 𝑆̅ is the averaged {110} surface area swept by the dislocation and 𝑉 is the volume of the sample. From equation (1), the amount of shear produced by a dislocation decreases as the size of the nanocube increases. As a consequence, and assuming a constant strain rate, the stress response to a shear increment will be smoother for large samples than for small one.
The two deformation regimes described above and their respective transition have already been observed in metallic nano-objects [10,90], and apply thus also to B1-structured ceramics. Further simulations e.g., discrete dislocation dynamics simulations (DDD) [65,91,92], performed at lower strain rates and applied to larger nanocubes should improve our understanding of the effective deformation processes that operate during in situ TEM compression tests of MgO nanocubes.
[27,76,79,86], we believe that the lattice friction is overcome in both the experiments and in the simulations of MgO nanocube leading to more isotropic dislocation lines.
This topological property is of primary importance because contact reactions between straight screw dislocations may sometimes not lead to the formation of junction locks. This particularly applies in the case of the BCC structure [93] and for the B1 structure of MgO as shown in recent DDD simulations [39]. Actually, non-screw dislocations are required to generate dislocation junction in MgO. Figure 4 shows images of a 12.6 nm nanocube deformed in MD simulations where two curved dislocations with 60 degrees tilted Burgers vectors react to create a junction embryo following the reaction path 1 2⁄ [101&](101) + 1 2⁄ [011&](011) → 1 2⁄ [112&](1&10). This local interaction process is not observed in smaller samples where all the deformation is accommodated by only few dislocations that never intersect (figure 4). This result confirms that the probability for dislocation to intersect is more likely in larger samples due to a higher number of defects. Finally, one can see figure 4 that contact interactions between dislocations (e.g., junctions, crossed or repulsive states) enhance the build-up of a dislocation network during the deformation of the 12.6 nm nanocube.
IV.4 Implications for NC ceramics
The results presented above have direct potential future impact for the processing (i.e. compaction, ball milling) or the design (i.e. mechanical properties) of nanostructured bulk ceramics. Indeed, knowing the plastic deformation mechanisms and mechanical constitutive laws of ceramic nanoparticles are of key importance for phenomena such as third body wear particle in contact mechanics, milling of nanoparticles which may exhibit a plastic behaviour below a certain size, or particle compaction during green body preparation of ceramics. It is for example usually considered that ceramic nanopowders just rearrange without plasticization during compaction [94]. The results presented here for MgO nanocubes and previously for transition alumina nanospheres [14,21] prove that wide plastic flow of ceramic nanoparticles can occur during compaction at RT. This propensity to plastic deformation during compaction opens a new route to deformable ceramics, even at RT. In other words, it would be possible to use plastic deformation of ceramic nanoparticles to obtain green bodies with very high densities and small pores, and sinter them at temperatures well below the ones currently used [95]. Sintering at much lower temperatures would then keep the nanoscale specificities after all the process chain, leading to higher mechanical or peculiar functional properties. This might then apply to orthopaedic devices, with ceramics exhibiting high wear and crack resistance, but also for transparent polycrystalline ceramics for which the quest is today to reach the highest density with the smallest grains [96].
V Conclusion
Ackowledgements
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